Ni-Cr-Co alloy for advanced gas turbine engines

ABSTRACT

A wrought age-hardenable nickel-chromium-cobalt based alloy suitable for use in high temperature gas turbine transition ducts possessing a combination of three specific key properties, namely resistance to strain age cracking, good thermal stability, and good creep-rupture strength contains in weight percent 17 to 22 chromium, 8 to 15 cobalt, 4.0 to 9.5 molybdenum, up to 7 tungsten, 1.28 to 1.65 aluminum, 1.50 to 2.30 titanium, up to 0.80 niobium, 0.01 to 0.2 carbon, up to 0.01 boron, and up to 3 iron, with a balance of nickel and impurities. Certain alloying elements must be present in amounts according to two equations here disclosed.

FIELD OF THE INVENTION

This invention relates to wroughtable high strength alloys for use atelevated temperatures. In particular, it is related to alloys whichpossess sufficient creep strength, thermal stability, and resistance tostrain age cracking to allow for fabrication and service in gas turbinetransition ducts and other gas turbine components.

BACKGROUND OF THE INVENTION

To meet the demand for increased operating efficiency, gas turbineengine designers would like to employ higher and higher operatingtemperatures. However, the ability to increase operating temperatures isoften limited by material properties. One application with such alimitation is gas turbine transition ducts. Transition ducts are oftenwelded components made of sheet or thin plate material and thus need tobe weldable as well as wroughtable. Often gamma-prime strengthenedalloys are used in transition ducts due to their high-strength atelevated temperatures. However, current commercially available wroughtgamma-prime strengthened alloys either do not have the strength orstability to be used at the very high temperatures demanded by advancedgas turbine design concepts, or can present difficulties duringfabrication. In particular, one such fabrication difficulty is thesusceptibility of many wrought gamma-prime strengthened alloys to strainage cracking. The problem of strain age cracking will be described inmore detail later in this document.

Wrought gamma-prime strengthened alloys are often based on thenickel-chromium-cobalt system, although other base systems are alsoused. These alloys will typically have aluminum and titanium additionswhich are responsible for the formation of the gamma-prime phase,Ni₃(Al,Ti). Other gamma-prime forming elements, such as niobium and/ortantalum, can also be employed. An age-hardening heat treatment is usedto develop the gamma-prime phase into the alloy microstructure. Thisheat treatment is normally given to the alloy when it is in the annealedcondition. The presence of gamma-prime phase leads to a considerablestrengthening of the alloy over a broad temperature range. Otherelemental additions may include molybdenum or tungsten for solidsolution strengthening, carbon for carbide formation, and boron forimproved high temperature ductility.

Strain age cracking is a problem which limits the weldability of manygamma-prime strengthened alloys. This phenomenon typically occurs when awelded part is subjected to a high temperature for the first time afterthe welding operation. Often this is during the post-weld annealingtreatment given to most welded gamma-prime alloy fabrications. Thecracking occurs as a result of the formation of the gamma-prime phaseduring the heat up to the annealing temperature. The formation of thestrengthening gamma-prime phase in conjunction with the low ductilitymany of these alloys possess at intermediate temperatures, as well asthe mechanical restraint typically imposed by the welding operation willoften lead to cracking. The problem of strain age cracking can limitalloys to be used up to only a certain thickness since greater materialthickness leads to greater mechanical restraint.

Several types of tests to evaluate the susceptibility of an alloy tostrain age cracking have been developed. These include the circularpatch test, the restrained plate test, and various dynamicthermal-mechanical tests. One test which can be used to evaluate thesusceptibility of an alloy to strain age cracking is the controlledheating rate tensile (CHRT) test developed in the 1960's. Recent testingat Haynes International has found the CHRT test to successfully rank thesusceptibility of several commercial alloys in an order consistent withfield experience. In the CHRT test, a sheet tensile sample is heatedfrom a low temperature up to the test temperature at a constant rate (arate of 25° F. to 30° F. per minute was used in the tests run at HaynesInternational). Once reaching the test temperature the sample is pulledto fracture at a constant engineering strain rate. The test samplestarts in the annealed (not age-hardened) condition, so that thegamma-prime phase is precipitating during the heat-up stage as would bethe case in a welded component being subjected to a post-weld heattreatment. The percent elongation to fracture in the test sample istaken as a measure of susceptibility to strain age cracking (lowerelongation values suggesting greater susceptibility to strain agecracking). The elongation in the CHRT is a function of test temperatureand normally will exhibit a minimum at a particular temperature. Thetemperature at which this occurs is around 1500° F. for many wroughtgamma-prime strengthened alloys.

Good strength and thermal stability at the high temperatures demanded byadvanced gas turbine concepts are two properties lacking in many currentcommercially available wrought gamma-prime strengthened alloys. Hightemperature strength has long been evaluated with the use ofcreep-rupture tests, where samples are isothermally subjected to aconstant load until the sample fractures. The time to fracture, orrupture life, is then used as a measure of the alloy strength at thattemperature. Thermal stability is a measure of whether the alloymicrostructure remains relatively unaffected during a thermal exposure.Many high-temperature alloys can form brittle intermetallic or carbidephases during thermal exposure. The presence of these phases candramatically reduce the room-temperature ductility of the material. Thisloss of ductility can be effectively measured using a standard tensiletest.

Many wrought gamma-prime strengthened alloys are available in sheet formtoday in today's marketplace. The Rene-41 or R-41 alloy (U.S. Pat. No.2,945,758) was developed by General Electric in the 1950's for use inturbine engines. It has excellent creep strength, but is limited by poorthermal stability and resistance to strain age cracking. A similarGeneral Electric alloy, M-252 alloy (U.S. Pat. No. 2,747,993), was alsodeveloped in the 1950's. Although currently available only in bar form,the composition would easily lend itself to sheet manufacture. The M-252alloy has good creep strength and resistance to strain age cracking, butlike R-41 alloy is limited by poor thermal stability. The Pratt &Whitney developed alloy known commercially as WASPALOY alloy (apparentlyhaving no U.S. patent coverage) is another gamma-prime strengthenedalloy intended for use in turbine engines and available in sheet form.However, this alloy has marginal creep strength above 1500° F., marginalthermal stability, and has fairly poor resistance to strain agecracking. The alloy commercially known as 263 alloy (U.S. Pat. No.3,222,165) was developed in the late 1950's and introduced in 1960 byRolls-Royce Limited. This alloy has excellent thermal stability andresistance to strain age cracking, but has very poor creep strength attemperatures greater than 1500° F. The PK-33 alloy (U.S. Pat. No.3,248,213) was developed by the International Nickel Company andintroduced in 1961. This alloy has good thermal stability and creepstrength, but is limited by a poor resistance to strain age cracking. Assuggested by these examples, no currently commercially available alloysare available which possess the unique combination of three keyproperties: good creep strength and good thermal stability in the 1600to 1700° F. temperature range as well as good resistance to strain agecracking.

SUMMARY OF THE INVENTION

The principal objective of this invention is to provide new wroughtage-hardenable nickel-chromium-cobalt based alloys which are suitablefor use in high temperature gas turbine transition ducts and other gasturbine components possessing a combination of three specific keyproperties, namely resistance to strain age cracking, good thermalstability, and good creep-rupture strength.

It has been found that this objective can be reached with an alloycontaining a certain range of chromium and cobalt, a certain range ofmolybdenum and possibly tungsten, and a certain range of aluminum,titanium and possibly niobium, with a balance of nickel and variousminor elements and impurities.

Specifically, the preferred ranges are 17 to 22 wt. % chromium, 8 to 15wt. % cobalt, 4.0 to 9.5 wt. % molybdenum, up to 7.0 wt. % tungsten,1.28 to 1.65 wt. % aluminum, 1.50 to 2.30 wt. % titanium, up to 0.80 wt.% niobium, up to 3 wt. % iron, 0.01 to 0.2 wt. % carbon, and up to 0.015wt. % boron, with a balance of nickel and impurities.

DESCRIPTION OF THE FIGURES

FIG. 1 is a graph of the ductility of the studied wrought age-hardenablenickel-chromium-cobalt based alloys in a controlled heating rate tensiletest at 1500° F.

FIG. 2 is a graph of the ductility of the studied wrought age-hardenablenickel-chromium-cobalt based alloys in a standard tensile test at roomtemperature.

DESCRIPTION OF THE PREFERRED EMBODIMENT

The wrought age-hardenable nickel-chromium-cobalt based alloys describedhere have sufficient creep strength, thermal stability, and resistanceto strain age cracking to allow for service in sheet or plate form ingas turbine transition ducts as well as in other product forms and otherdemanding gas turbine applications. This combination of criticalproperties is achieved through control of several critical elements eachwith certain functions. The presence of gamma-prime forming elementssuch as aluminum, titanium, and niobium contribute significantly to thehigh creep-rupture strength through the formation of the gamma-primephase during the age-hardening process. However, the combined amount ofaluminum, titanium, and niobium must be carefully controlled to allowfor good resistance to strain age cracking. Molybdenum and possiblytungsten are added to provide additional creep-rupture strength throughsolid solution strengthening. Again, however, the total combinedmolybdenum and tungsten concentration must be carefully controlled, inthis case to ensure sufficient thermal stability of the alloy.

Based on the projected requirements for the next generation of gasturbine transition ducts, gamma-prime strengthened alloys havesignificant potential. Three of the more critical properties are creepstrength, weldability (i.e. strain age cracking resistance), and thermalstability. However, producing a gamma-prime strengthened alloy whichexcels in all three of these properties is not straightforward and nocommercially available alloy was found which possessed all threeproperties to a sufficient degree.

I tested 26 experimental and 5 commercial alloys whose compositions areset forth in Table 1. The experimental alloys have been labeled Athrough Z. The commercial alloys were HAYNES R-41 alloy, HAYNES WASPALOYalloy, HAYNES 263 alloy, M-252 alloy, and NIMONIC PK-33 alloy. Thealloys (including both the experimental and the commercial alloys) had aCr content which ranged from 17.5 to 21.3 wt. %, as well as a cobaltcontent ranging from 8.3 to 19.6 wt. %. The aluminum content ranged from0.49 to 1.89 wt. %, the titanium content from 1.53 to 3.12 wt. %, andthe niobium content ranged from nil to 0.79 wt. %. The molybdenumcontent ranged from 3.2 to 10.5 wt. % and the tungsten ranged from nilup to 8.3 wt. %. Intentional minor element additions carbon and boronranged from 0.034 to 0.163 wt. % and from nil to 0.008 wt. %,respectively. Iron ranged from nil to 3.6 wt. %.

All testing of the alloys was performed on sheet material of 0.047″ to0.065″ thickness. The experimental alloys were vacuum induction melted,and then electro-slag remelted, at a heat size of 50 lb. The ingots soproduced were soaked at 2150° F. and then forged and rolled withstarting temperatures of 2150° F. The sheet thickness after hot rollingwas 0.085″. The sheets were annealed at 2150° F. for 15 minutes andwater quenched. The sheets were then cold rolled to 0.060″ thickness.The cold rolled sheets were annealed at temperatures between 2050 and2175° F. as necessary to produce a fully recrystallized, equiaxed grainstructure with an ASTM grain size between 4 and 5. Finally, the sheetmaterial was given an age-hardening heat treatment of 1475° F. for 8hours to produce the gamma-prime phase. The commercial alloys HAYNESR-41 alloy, HAYNES WASPALOY alloy, HAYNES 263 alloy, and NIMONIC PK-33alloy were obtained in sheet form in the mill annealed condition. Sinceno commercially available M-252 alloy sheet could be found, a 50 lb.heat was produced for evaluation using the same method as describedabove for the experimental alloys. All five of the commercial alloyswere given post-anneal age-hardening heat treatments according acceptedstandards. These heat treatments are reported in Table 2.

To evaluate the three properties identified above as important (strainage cracking resistance, thermal stability, and creep strength) threedifferent tests were employed on each of the alloys. The first test wasthe controlled heating rate tensile test (CHRT). The results of the CHRTtesting are given in Table 3. The critical property in this test is thetensile ductility, as measured by a measurement of the elongation tofailure. Alloys with a greater ductility in this test are expected tohave greater resistance to strain age cracking. The objective of thepresent study was to have a ductility of 4.5% or greater. Of theexperimental alloys, only alloy W failed to meet this requirement. Forthe commercial alloys, M-252 alloy and 263 alloy met the requirement,while PK-33 alloy, WASPALOY alloy, and R-41 alloy did not. It was foundthat the performance of a given alloy in the CHRT test could becorrelated to the amount of the gamma-prime forming elements in thealloy using the following equation (where the elemental compositions arein wt. %):Al+0.56Ti+0.29Nb<2.9  (1)

The values of the left hand side of Eq. (1) for all of the alloys inthis study are given in Table 1. All of the alloys which passed the CHRTtest were found to obey Eq. (1). Furthermore, all of the alloys whichdid not obey Eq. (1) did not pass the CHRT test requirement, that is,they were found to have a 1500° F. CHRT ductility less than 4.5%. Thisrelationship is shown more clearly in FIG. 1, where the 1500° F. CHRTductility is plotted against the value of the left hand side of Eq. (1)for all of the alloys in the study. All testing was performed on samplesin the annealed condition. The tensile ductility (measured as thepercent elongation to failure) is plotted as a function of thecompositional variable Al+0.56Ti+0.29Nb (where the elementalcompositions are in wt. %). A line is drawn on the figure correspondingto a tensile ductility of 4.5%. All alloys plotted above this line(symbol: filled circles) were considered to have passed the controlledheating rate tensile test, while alloys plotted below the line (symbol:x-marks) were considered to have failed. A dashed vertical line is drawnat a value of 2.9 wt. % for the compositional variable,Al+0.56Ti+0.29Nb. All alloys with a value greater than 2.9 were found tofail the controlled heating rate tensile test. TABLE 1 Alloy Ni Cr Co WMo Ti Al Nb C B Fe Mo + 0.52W Al + 0.56Ti + 0.29Cb A BAL 19.1 10.7 7.05.5 1.91 1.53 <0.05 0.079 0.003 <0.1 9.1 2.60 B BAL 19.5 10.9 5.4 4.32.07 1.51 0.02 0.097 0.006 <0.1 7.1 2.68 C BAL 19.2 10.8 6.3 5.1 2.201.60 <0.05 0.095 0.006 <0.1 8.4 2.84 D BAL 19.0 10.7 5.9 6.3 1.71 1.570.63 0.090 0.005 <0.1 9.3 2.71 E BAL 19.2 10.7 6.8 5.3 1.59 1.51 0.790.085 0.003 <0.1 8.8 2.63 F BAL 19.3 10.8 6.1 4.8 2.08 1.51 <0.05 0.0910.002 <0.1 8.0 2.68 G BAL 19.1 10.7 6.3 5.1 2.03 1.40 <0.05 0.097 0.002<0.1 8.4 2.54 H BAL 19.3 10.8 6.1 4.6 1.85 1.63 <0.05 0.088 0.003 0.27.8 2.67 I BAL 19.3 10.7 6.1 4.7 1.89 1.29 <0.05 0.075 0.004 0.2 7.92.35 J BAL 19.2 10.7 6.1 4.6 2.28 1.30 <0.05 0.074 0.003 0.2 7.8 2.58 KBAL 20.3 10.8 <0.1 7.8 2.06 1.51 <0.05 0.065 0.006 0.2 7.8 2.67 L BAL19.2 10.8 6.0 4.8 2.08 1.48 0.02 0.088 0.005 <0.1 7.9 2.65 M BAL 19.310.7 6.1 4.6 1.97 1.39 <0.05 0.081 0.003 2.6 7.8 2.50 N BAL 21.3 8.3 6.04.7 2.13 1.45 <0.05 0.073 0.004 0.2 7.8 2.65 O BAL 17.5 14.2 6.1 4.72.11 1.47 <0.05 0.077 0.004 0.2 7.9 2.66 P BAL 19.4 10.7 6.2 4.6 1.981.52 <0.05 0.034 0.006 0.2 7.8 2.64 Q BAL 19.2 10.7 2.7 6.2 2.01 1.54<0.05 0.056 0.006 0.2 7.6 2.68 R BAL 19.9 10.1 <0.1 7.2 2.05 1.50 <0.050.058 0.006 0.7 7.2 2.65 S BAL 20.2 9.6 <0.1 8.3 2.12 1.48 <0.05 0.0620.007 0.7 8.3 2.67 T BAL 18.9 10.1 <0.1 9.3 2.07 1.56 <0.05 0.066 0.0060.7 9.3 2.72 U BAL 18.7 10.5 8.3 6.3 1.80 1.43 <0.05 0.089 0.002 0.110.6 2.44 V BAL 19.6 10.9 0.1 9.9 2.21 1.33 0.65 0.094 0.004 <0.1 9.92.76 W BAL 19.4 10.9 5.4 4.3 2.30 1.66 <0.05 0.096 0.006 <0.1 7.1 2.95 XBAL 18.8 10.3 7.6 6.0 1.53 1.39 0.72 0.089 <0.002 <0.1 9.9 2.46 Y BAL19.2 10.6 4.1 3.2 2.13 1.45 <0.05 0.080 0.004 0.2 5.3 2.65 Z BAL 19.210.8 0.1 10.5 2.10 1.46 <0.05 0.077 0.004 0.2 10.5 2.64 M-252 BAL 18.99.7 <0.1 10.0 2.30 1.01 0.04 0.163 0.005 0.2 10.0 2.31 PK-33 BAL 18.813.1 — 7.2 1.90 1.89 — 0.048 0.003 0.7 7.2 2.95 263 BAL 20.5 19.6 <0.15.9 2.16 0.49 <0.05 0.060 0.002 0.4 5.9 1.61 WASP BAL 19.1 13.3 <0.1 4.32.92 1.45 0.05 0.080 0.008 1.0 4.3 2.97 R-41 BAL 19.1 10.9 <0.1 9.7 3.121.48 <0.05 0.090 0.008 3.6 9.7 3.10

TABLE 2 Alloy Heat Treatments* Experimental alloys A-Z 1475° F./8 hr./ACR-41 alloy 2050° F./30 min./AC + 1650° F./4 hr./AC WASPALOY alloy 1825°F./2 hr./AC + 1550° F./4 hr./AC + 1400° F./16 hr./AC 263 alloy 1472°F./8 hr./AC M-252 alloy 1400° F./15 hr./AC PK-33 alloy 1562° F./4 hr./AC*All heat treatments performed after an annealing heat treatment.AC = air cool

TABLE 3 Alloy 1500° F. CHRT Ductility (% Elong.) A 5.9 B 4.9 C 5.0 D 6.4E 9.5 F 6.1 G 4.9 H 8.5 I 10.0 J 5.5 K 8.3 L 5.7 M 8.5 N 5.6 O 5.8 P 5.2Q 5.9 R 6.9 S 8.2 T 7.0 U 5.0 V 6.7 W 4.2 X 6.9 Y 5.1 Z 9.3 R-41 alloy2.8 WASPALOY alloy 3.5 263 alloy 22.9 M-252 alloy 5.6 PK-33 alloy 3.6

To evaluate the thermal stability of the alloys, their room temperaturetensile ductility was determined after a long term thermal exposure.After performing the age-hardening heat treatments given in Table 2,samples from all of the experimental and commercial alloys were given athermal exposure of 1600° F./1000 hrs./AC. A room temperature tensiletest was performed on the thermally exposed samples and the results aregiven in Table 4. Ductility greater than 20% was considered acceptable.Using this guideline, the experimental alloys U, V, X, and Z were foundto fail along with the commercial alloys M-252 alloy, WASPALOY alloy,and R-41 alloy. It was found that control of the elements molybdenum andtungsten was critical to develop a thermally stable alloy. The followingrelationship was found (where the elemental compositions are in wt. %):Mo+0.52W<9.5  (2)

The values of the left hand side of Eq. (2) for all of the alloys inthis study are given in Table 1. All of the alloys which did not obeyEq. (2) were found to not have sufficient thermal stability, that is,their room temperature tensile ductility after a 1000 hour thermalexposure at 1600° F. was found to be less than 20%. One alloy (WASPALOYalloy) was found to satisfy Eq. (2), but to have poor thermal stability.However, this alloy did not satisfy Eq. (1) and therefore is notsuitable for the target application. From this example, it is clear thatto ensure thermal stability for this class of alloys, it is necessary tocontrol the amount of aluminum, titanium, and niobium as well as themolybdenum and tungsten. The usefulness of Eq. (2) becomes quite clearwhen considering FIG. 2, where the ductility of the thermally exposedsamples is plotted against the value of the left hand side of Eq. (2)for all of the alloys in the study. Only alloys which satisfy therelationship Al+0.56Ti+0.29Nb<2.9 (where the elemental compositions arein wt. %) are plotted in the graph. All testing was performed on samplesgiven an age-hardening heat treatment followed by a thermal exposure of1600° F. for 1000 hours. In the graph, the tensile ductility (measuredas the percent elongation to failure) is plotted as a function of thecompositional variable Mo+0.52W (where the elemental compositions are inwt. %). A line is drawn on the figure corresponding to a tensileductility of 20%. All alloys plotted above this line (symbol: filledcircles) were considered to have passed the thermal stability test,while alloys plotted below the line (symbol: x-marks) were considered tohave failed. A dashed vertical line is drawn at a value of 9.5 wt. % forthe compositional variable, Mo+0.52W. All alloys with a value greaterthan 9.5 were found to fail the thermal stability test. TABLE 4 AlloyDuctility after 1600° F./1000 hrs./AC (% Elong.) A 27.8 B 29.2 C 28.8 D22.2 E 24.3 F 28.2 G 26.3 H 29.3 I 34.3 J 30.8 K 31.4 L 30.2 M 32.1 N23.5 O 32.5 P 32.8 Q 29.4 R 34.5 S 33.6 T 29.9 U 10.4 V 9.2 W 27.3 X19.0 Y 33.6 Z 18.0 R-41 alloy 2.6 WASPALOY 12.8 263 alloy 40.9 M-252alloy 10.1 PK-33 alloy 26.2

The third key property for the target application is creep strength. Thecreep-rupture strength of the alloys was measured at 1700° F. with aload of 7 ksi. A rupture life of greater than 300 hours was theestablished goal. The results for the experimental and commercial alloysare shown in Table 5. All of the experimental alloys were found to passthe goal, with the exception of alloys V, Y, and Z. The commercialalloys all passed with the exception of 263 alloy and WASPALOY alloy. Ofthe total of five alloys which failed the creep-rupture goal, three ofthem (alloys V and Z, as well as WASPALOY alloy) did not satisfy one orboth of Eqs. (1) and (2) and were thermally unstable. Thermalinstability can be a negative influence on creep strength. The other twoalloys which did not meet the creep strength goal (alloy Y and 263alloy) both had a relatively low total content of the solid solutionstrengthening elements molybdenum and tungsten. Additionally, the 263alloy had a low total content of the gamma-prime forming elementsaluminum, titanium, and niobium. To ensure adequate levels of both thesolid solution strengthening elements and the gamma-prime formingelements, the Eqs. (1) and (2) were modified respectfully as (where theelemental compositions are in wt. %):2.2<Al+0.56Ti+0.29Nb<2.9  (3)and6.5<Mo+0.52W<9.5  (4)

Of the 31 total experimental and commercial alloys tested in this study,20 were found to pass all three key property tests, i.e. the CHRT test,the thermal exposure test, and the creep-rupture test. All 20 of theacceptable alloys (experimental alloys A through T) had compositionswhich satisfied both Eqs. (3) and (4). The 11 alloys which were deemedunacceptable (which included experimental alloys U through Z and allfive of the commercial alloys) had compositions which failed to satisfyone or both of Eqs. (3) and (4). From Table 1 it can be seen that theacceptable alloys contained in weight percent 17.5 to 21.3 chromium, 8.3to 14.2 cobalt, 4.3 to 9.3 molybdenum, up to 7.0 tungsten, 1.29 to 1.63aluminum, 1.59 to 2.28 titanium, up to 0.79 niobium, 0.034 to 0.097carbon, 0.002 to 0.007 boron and up to 2.6 iron. For the reasonsexplained below, alloys containing these elements within the followingranges and meeting Eqs. (3) and (4) should provide the desiredproperties: 17 to 22 chromium, 8 to 15 cobalt, 4.0 to 9.5 molybdenum, upto 7.0 tungsten, 1.28 to 1.65 aluminum, 1.50 to 2.30 titanium, up to0.80 niobium, 0.01 to 0.2 carbon and up to 0.015 boron with the balancebeing nickel plus impurities. The alloy may also contain tantalum, up to1.5 wt. %, manganese, up to 1.5 wt. %, silicon, up to 0.5 wt. %, and oneor more of magnesium, calcium, hafnium, zirconium, yttrium, cerium andlanthanum. Each of these seven elements may be present up to 0.05 wt. %.The acceptable alloys had a range of values for Al+0.56 Ti+0.29 Nb offrom 2.35 to 2.84 and a range for Mo+0.52W of from 7.1 to 9.3. TABLE 5Alloy Rupture Life (hours) A 304 B 560 C 481 D 375 E 346 F 509 G 584 H764 I 410 J 767 K 460 L 522 M 581 N 401 O 403 P 664 Q 419 R 328 S 641 T506 U 384 V 284 W 463 X 339 Y 271 Z 283 R-41 alloy 618 WASPALOY 243 263alloy 139 M-252 alloy 392 PK-33 alloy 412

The presence of chromium (Cr) in alloys used in high temperatureenvironments provides for necessary oxidation and hot corrosionresistance. In general, the higher the Cr content the better theoxidation resistance, however, too much Cr can lead to thermalinstability in the alloy. For the alloys of this invention, it was foundthat the chromium level should be between about 17 to 22 wt. %.

Cobalt (Co) is a common element in many wrought gamma-prime strengthenedalloys. Cobalt decreases the solubility of aluminum and titanium innickel at lower temperatures allowing for a greater gamma-prime contentfor a given level of aluminum and titanium. It was found that Co levelsof about 8 to 15 wt. % are acceptable for the alloys of this invention.

As mentioned previously, aluminum (Al), titanium (Ti), and niobium (Nb)contribute to the creep-strength of the alloys of this invention throughthe formation of the strengthening gamma-prime phase upon anage-hardening heat treatment. The combined total of these elements islimited by Eq. (3) above. In terms of the individual elements, it wasfound that Al could range from 1.28 to 1.65 wt. %, Ti could range from1.50 to 2.30 wt. %, and Nb could range from nil to 0.80 wt. %.

As mentioned previously, molybdenum (Mo) and tungsten (W) contribute tothe creep-rupture strength of the alloys of this invention through solidsolution strengthening. The combined total of these elements is limitedby Eq. (4) above. In terms of the individual elements, it was found thatMo could range from about 4.0 to 9.5 wt. %, while W could range from nilto about 7.0 wt. %.

Carbon (C) is a necessary component and contributes to creep-strength ofthe alloys of this invention through formation of carbides. Carbides arealso necessary for proper grain size control. Carbon should be presentin the amount of about 0.01 to 0.2 wt. %.

Iron (Fe) is not required, but typically will be present. The presenceof Fe allows economic use of revert materials, most of which containresidual amounts of Fe. An acceptable, Fe-free alloy might be possibleusing new furnace linings and high purity charge materials. Thepresented data indicate that levels up to at least about 3 wt. % areacceptable.

Boron (B) is normally added to wrought gamma-prime strengthened alloysin small amounts to improve elevated temperature ductility. Too muchboron may lead to weldability problems. The preferred range is up toabout 0.015 wt. %.

Tantalum (Ta) is a gamma-prime forming element in this class of alloys.It is expected that tantalum could be partially substituted foraluminum, titanium, or niobium at levels up to about 1.5 wt. %.

Manganese (Mn) is often added to nickel based alloys to help controlproblems arising from the presence of sulfur impurities. It is expectedthat Mn could be added to alloys of this invention to levels of at least1.5 wt. %.

Silicon (Si) can be present as an impurity and is sometimesintentionally added for increased environmental resistance. It isexpected that Si could be added to alloys of this invention to levels ofat least 0.5 wt. %.

Copper (Cu) can be present as an impurity originating either from theuse of revert materials or during the melting and processing of thealloy itself. It is expected that Cu could be present in amounts up toat least 0.5 wt. %.

The use of magnesium (Mg) and calcium (Ca) is often employed duringprimary melting of nickel base alloys. It is expected that levels ofthese elements up to about 0.05 wt. % could be present in alloys of thisinvention.

Often, small amounts of certain elements are added to nickel basedalloys to provide increased environmental resistance. These elementsinclude, but are not necessarily limited to lanthanum (La), cerium (Ce),yttrium (Y), zirconium (Zr), and hafnium (Hf). It is expected thatamounts of each of these elements up to about 0.05 wt. % could bepresent in alloys of this invention.

Even though the samples tested were limited to wrought sheet, the alloysshould exhibit comparable properties in other wrought forms (such asplates, bars, tubes, pipes, forgings, and wires) and in cast,spray-formed, or powder metallurgy forms, namely, powder, compactedpowder and sintered compacted powder. Consequently, the presentinvention encompasses all forms of the alloy composition.

The combined properties of good thermal stability, resistance to strainage cracking and good creep rupture strength exhibited by this alloymake it particularly useful for fabrication into gas turbine enginecomponents and particularly useful for transition ducts in theseengines. Such components and engines containing these components can beoperated at higher temperatures without failure and should have a longerservice life than those components and engines currently available.

Although I have disclosed certain preferred embodiments of the alloy, itshould be distinctly understood that the present invention is notlimited thereto, but may be variously embodied within the scope of thefollowing claims.

1. A nickel-chromium-cobalt based alloy having a composition comprisedin weight percent of: 17 to 22 chromium 8 to 15 cobalt 4.0 to 9.5molybdenum up to 7.0 tungsten 1.28 to 1.65 aluminum 1.50 to 2.30titanium up to 0.80 niobium 0.01 to 0.2 carbon up to 0.015 boron with abalance of nickel and impurities, the alloy further satisfying thefollowing compositional relationships defined with elemental quantitiesbeing in terms of weight percent:2.2<Al+0.56Ti+0.29Nb<2.96.5<Mo+0.52W<9.5
 2. The nickel-chromium-cobalt based alloy of claim 1,also containing iron up to 3 weight percent.
 3. Thenickel-chromium-cobalt based alloy of claim 1, also containing in weightpercent at least one of tantalum, up to 1.5%, manganese, up to 1.5%,silicon, up to 0.5%, and copper, up to 0.5%.
 4. Thenickel-chromium-cobalt based alloy of claim 1, also containing at leastone element selected from the group consisting of magnesium, calcium,hafnium, zirconium, yttrium, cerium, and lanthanum, wherein each saidelement present comprises up to 0.5 weight percent of the alloy.
 5. Thenickel-chromium-cobalt based alloy of claim 1, wherein the alloy is inwrought form selected from the group consisting of sheets, plates, bars,wires, tubes, pipes, and forgings.
 6. The nickel-chromium-cobalt basedalloy of claim 1, wherein the alloy is in cast form.
 7. Thenickel-chromium-cobalt based alloy of claim 1, wherein the alloy hasbeen spray-formed.
 8. The nickel-chromium-cobalt based alloy of claim 1,wherein the alloy is in powder metallurgy form.
 9. Thenickel-chromium-cobalt based alloy of claim 1 wherein the alloy isformed as a component for a gas turbine engine.
 10. Anickel-chromium-cobalt based alloy, suitable for use in gas turbinetransition ducts, having a composition in weight percent consistingessentially of: 17.5 to 21.3 chromium 8.3 to 14.2 cobalt 4.3 to 9.3molybdenum up to 7.0 tungsten 1.29 to 1.63 aluminum 1.59 to 2.28titanium up to 0.79 niobium 0.034 to 0.097 carbon 0.002 to 0.007 boronup to 2.6 iron with a balance of nickel and impurities, the alloyfurther satisfying the following compositional relationships definedwith elemental quantities being in terms of weight percent:2.35<Al+0.56Ti+0.29Nb<2.847.1<Mo+0.52W<9.3
 11. The nickel-chromium-cobalt based alloy of claim 10,also containing in weight percent at least one of tantalum, up to 1.5%,manganese, up to 1.5%, silicon, up to 0.5%, and copper, up to 0.5%. 12.The nickel-chromium-cobalt based alloy of claim 10, also containing upto at least one element selected from the group consisting of magnesium,calcium, hafnium, zirconium, yttrium, cerium, and lanthanum, whereineach said element present comprises up to 0.05 weight percent of thealloy.
 13. The nickel-chromium-cobalt based alloy of claim 10, whereinthe alloy is in wrought forms selected from the group consisting ofsheets, plates, bars, wires, tubes, pipes, and forgings.
 14. Thenickel-chromium-cobalt based alloy of claim 10, wherein the alloy is incast form.
 15. The nickel-chromium-cobalt based alloy of claim 10,wherein the alloy has been spray-formed.
 16. The nickel-chromium-cobaltbased alloy of claim 10, wherein the alloy is in powder metallurgy form.17. The nickel-chromium-cobalt based alloy of claim 10 wherein the alloyis formed as a component for a gas turbine engine.
 18. An improved gasturbine engine of the type having a plurality of metal componentswherein the improvement comprises at least one of the metal componentsis comprised in weight percent of: 17 to 22 chromium 8 to 15 cobalt 4.0to 9.5 molybdenum up to 7.0 tungsten 1.28 to 1.65 aluminum 1.50 to 2.30titanium up to 0.80 niobium 0.01 to 0.2 carbon up to 0.015 boron with abalance of nickel and impurities, the alloy further satisfying thefollowing compositional relationships defined with elemental quantitiesbeing in terms of weight percent:2.2<Al+0.56Ti+0.29Nb<2.96.5<Mo+0.52W<9.5
 19. The improved gas turbine engine wherein the atleast one of the metal components is a transition duct.
 20. The improvedgas turbine engine of claim 18 where the at least one of the metalcomponents in weight percent consists essentially of: 17.5 to 21.3chromium 8.3 to 14.2 cobalt 4.3 to 9.3 molybdenum up to 7.0 tungsten1.29 to 1.63 aluminum 1.59 to 2.28 titanium up to 0.79 niobium 0.034 to0.097 carbon 0.002 to 0.007 boron up to 2.6 iron with a balance ofnickel and impurities, the alloy further satisfying the followingcompositional relationships defined with elemental quantities being interms of weight percent:2.35<Al+0.56Ti+0.29Nb<2.847.1<Mo+0.52W<9.3